Method for manufacturing alloy containing transition metal carbide, tungsten alloy containing transition metal carbide, and alloy manufactured by said method

ABSTRACT

The present invention relates to the development of an alloy material with significantly improved low-temperature brittleness, recrystallization brittleness, and irradiation brittleness by the introduction of a recrystallization microstructure into an alloy, particularly a tungsten material, to significantly strengthen a weak grain boundary of the recrystallization microstructure. The present invention comprises the steps of: mechanically alloying at least one species selected from a group-IVA, VA, or VIA transition metal carbide and a metallic raw material; sintering base powders obtained through the mechanically alloying step, by using a hot isostatic press; and performing plastic deformation of at least 60% on the alloy obtained through the sintering step, at a strain rate between 10 −5  s −1  and 10 −2  S −1  and at a temperature between 500° C. and 2,000° C. It is therefore possible to obtain an alloy material with significantly improved low-temperature brittleness, recrystallization brittleness, and irradiation brittleness.

TECHNICAL FIELD

The present invention relates to a method for manufacturing an alloycontaining transition metal carbide, a tungsten alloy containingtransition metal carbide, and an alloy manufactured by said method. Inparticular, the present invention relates to a method for manufacturingan alloy that manifests superplasticity due to grain boundary slidingwhen the alloy is made to undergo superplastic deformation, thatexhibits high recrystallization fracture strength, that has littledecrease in strength or ductility, even when heated to high temperaturesdue to its recrystallized structure, and which has dramatically remediedlow-temperature embrittlement, recrystallization embrittlement, andneutron irradiation embrittlement, as well as an alloy that has beenmanufactured by this manufacturing method, in particular, a tungstenalloy.

BACKGROUND ART

Tungsten and tungsten alloys have melting points of as high as 3410° C.which are the highest of any metal. These materials thus provide a greatmany advantages that are unparalleled by other metals. However, thematerials have not been used for structure due to an inability toresolve problems with persisting embrittlement (low-temperatureembrittlement, recrystallization embrittlement, and irradiationembrittlement), which has hampered practical use of these materials ashigh-temperature structural materials in extreme environments.

These embrittlement phenomena all result from very weak crystal grainboundaries and a tendency for fracture to originate from the grainboundaries which are referred to as “grain boundary embrittlement”. Thecause of grain boundary embrittlement is that tungsten is a metal havingan extremely high degree of covalent bond character, and the grainboundaries are substantially weaker (tend to fracture) due to their highenergies. Additionally, interstitial gas elements contained in air suchas nitrogen and oxygen have extremely low solubility in tungsten, andthus tend to segregate and precipitate at the grain boundaries, whichfurther weakens the grain boundaries and promotes embrittlement.

As shown in FIG. 1( a), with common metals, almost the entiretemperature range is the ductile temperature region, because of the factthat plastic deformation (permanent set) occurs prior to break. On theother hand, as shown in FIG. 1( b), because tungsten has covalentbonding in which the directionality of the interatomic bonds isextremely high, the grain boundaries are substantially weaker,ductile-brittle transition occurs, and the ductile-brittle transitiontemperature (“DBTT” below) is also high. This phenomenon thus becomesextreme with lower temperatures at which there is a precipitous increasein the Peierls stress (yield strength) required for screw dislocationsto glide in tungsten (low-temperature embrittlement), and the phenomenonis even more pronounced with recrystallized structure in which extremelyweak grain boundaries are formed (recrystallization embrittlement).Moreover, when lattice defects are introduced by high-energy particleirradiation using neutrons or the like, such irradiation induced defectsaccumulate inside the crystal grains or at the grain boundaries andimpede dislocation slip, resulting in the promotion of grain boundaryembrittlement (irradiation embrittlement).

Consequently, in order to simultaneously remedy low-temperatureembrittlement, recrystallization embrittlement, and irradiationembrittlement, it is necessary to introduce recrystallizedmicrostructures containing high densities of sink sites (sinks; crystalgrain boundary or dispersed particles) that can permit the material totolerate irradiation induced defects and to convert the weak grainboundaries in the recrystallized microstructure to strong grainboundaries that resist fracture.

The inventors of the present invention, in order to resolve problemswith low temperature embrittlement and embrittlement of tungsten due toneutron irradiation and recrystallization, carried out manufacture ofW-TiC having ultrafine crystal grains by a hot isostatic pressing (HIP)method and by mechanical alloying (MA) in Ar and H₂ atmospheres. Effectssuch as an appreciable improvement in room temperature toughness werefound to occur, and these results were published (non-patent documents1, 2). However, the tungsten materials manufactured by the above methodsstill were not adequate for practical use.

On the other hand, a known method for improving toughness and the likeof high-melting (refractory) metals has been to increase creepresistance by the introduction of 0.005 to 10 mass % of one or moretypes of compounds or mixtures selected from the group consisting ofoxides, nitrides, carbides, borides, silicates, or aluminates withparticle diameters of ≦1.5 μm and having melting points of 1500° C. orgreater into high-melting metals such as Mo, W, Nb, Ta, V, and Cr (referto patent document 1). However, patent document 1 discloses animprovement in the heat resistance and creep resistance of high-meltingmetals at high temperatures, not a remedy for low-temperatureembrittlement, recrystallization embrittlement, or irradiationembrittlement.

In addition, the inventors of the present invention also applied for apatent (refer to patent document 2) based on the discovery thatdispersion of 0.05 to 5 mol % of ultrafine particles of group IVatransition metal carbides with particle diameters of 10 nm or less inmolybdenum alloy and restricting the crystal grain diameter to 1 μm orless enables the strength of the molybdenum alloy to increase, less lossof strength, even when heated at high-temperature, and an alleviation oflow-temperature embrittlement, recrystallization embrittlement, andneutron irradiation embrittlement. However, the molybdenum described inpatent document 2 is a material that exhibits ductility at roomtemperature, even as a pure metal, and has completely differentproperties and manufacture conditions in comparison to tungsten, whichis an extremely brittle material having a high melting point that is800° C. higher than that of molybdenum.

In addition, with molybdenum, it is necessary to introduce work-deformedstructure by plastic working (hammering (forging), rolling, or the like)in order to improve ductility in patent document 2, resulting in adecrease in recrystallization temperature and anisotropy. With tungsten,on the other hand, the issue is ductility improvement in arecrystallized equiaxed structure that is in a recrystallized state withabsolutely no work-deformed structure and therefore no anisotropy. Thetwo cases are thus substantially different.

PRIOR ART DOCUMENTS

Patent Document

-   Patent document 1: Japanese Laid-open Patent Publication No.    1-502680-   Patent document 2: Japanese Laid-open Patent Publication No. 8-85840

Non-Patent Documents

-   Non-patent document 1: Collected abstracts of the Japan Institute of    Metals and Materials Vol. 148, p. 235-   Non-patent document 2: Collected abstracts of the Japan Institute of    Metals and Materials Vol. 143, p. 322

DISCLOSURE OF THE INVENTION

Problems to be Solved by the Invention

The inventors of the present invention carried out painstakinginvestigations and discovered that, when an alloy containing transitionmetal carbide that has been produced by mechanical alloying (MA) and hotisostatic pressing (HIP) is additionally subjected to a strengtheningtreatment for the recrystallized random grain boundaries employing grainboundary sliding by superplastic deformation in order to reinforce therecrystallized grain boundaries in the alloy, the weak grain boundariesin the recrystallized microstructure can be dramatically strengthened.As a result, there is a dramatic resolution of low-temperatureembrittlement, recrystallization embrittlement, and irradiationembrittlement. In addition, although the strengthening treatment forrandom recrystallized grain boundaries employing grain boundary slidingby superplastic deformation can be applied to any alloy that exhibitssuperplastic deformation due to grain boundary sliding, the noveldiscovery was made that this method is effective for reducing thebrittleness of tungsten, which is the most brittle material among themetals. The present invention was realized based on this new knowledge.

Specifically, an aim of the present invention is to provide a method formanufacturing alloys containing transition metal carbide havingdramatically resolved low-temperature embrittlement, recrystallizationembrittlement, and irradiation embrittlement. An additional aim of thepresent invention is to provide an alloy that is manufactured by thismanufacturing method. An additional aim of the present invention is toprovide a tungsten alloy containing a transition metal carbide havingdramatically resolved low-temperature embrittlement, recrystallizationembrittlement, and irradiation embrittlement.

Means for Resolving the Problems

The present invention is described below and relates to a method formanufacturing an alloy containing a transition metal carbide, a tungstenalloy containing a transition metal carbide, and an alloy that ismanufactured by this manufacturing method.

(1) A method for manufacturing an alloy, characterized by having a stepfor mechanically alloying a metal raw material and at least one selectedfrom carbides of group IVA, VA, or VIA transition metals, a step forsintering the raw material powder obtained in the mechanical alloyingstep (i.e., mechanically alloyed powder) using hot isostatic pressing(HIP), and a step for subjecting the alloy obtained in the sinteringstep to superplastic deformation due to grain boundary sliding of 60% orgreater at 500 to 2000° C. and at a strain rate of 10⁻⁵ to 10⁻² s⁻¹.

(2) The method for manufacturing an alloy according to (1),characterized by having a step in which the transition metal carbide andthe metal raw material are degassed by heating prior to the mechanicalalloying step.

(3) A tungsten alloy comprising 0.25 to 5 mass % of at least one typeselected from carbides of a group IVA, VA, or VIA transition metals, thetungsten alloy characterized in that the oxygen content is 950 ppm bymass or less, the nitrogen content is 60 ppm by mass or less, 80% ormore of the tungsten phase observed in a sectioned surface area isrecrystallized to equiaxed grains with grain diameters of 0.05 to 10 μm,the ductile-brittle transition temperature determined by three-pointflexure is 500K or less, and plastic deformation is possible at or abovethis temperature.

(4) The tungsten alloy according to (3), characterized in that 90% orgreater of the azimuths of the carbide present in the tungsten alloystructure and the azimuths of the tungsten matrix are in the(Kurdjumov-Sachs) azimuth relationship: {111} W//{110} transition metalcarbide <110> W//<111> transition metal carbide.

(5) The tungsten alloy according to (3) or (4), characterized in thatthe full width at half maximum for reflection of the (220) diffractionplanes is 3° or less as determined by X-ray diffraction, or that thereare 50 or fewer dislocations within crystal grains as determined bytransmission electron microscopy.

(6) The tungsten alloy according to any of (3) to (5), characterized inthat the maximum bend strength determined by three-point flexure is 1470MPa or greater.

(7) The alloy that is manufactured by the manufacturing method accordingto (1) or (2).

Effect of the Invention

In accordance with the present invention, transition metal carbide andalloy powder are treated by mechanical alloying (MA) method and hotisostatic pressing (HIP) method, and superplastic deformation that canmaximally utilize grain boundary sliding is used in order to foster andoptimize carbide grain boundary precipitation and grain boundarysegregation in recrystallized micrograin structures. As a result, (1)the grain boundary strength of the alloy in the recrystallized structureis significantly improved, particularly the grain boundary strength(grain boundary bond strength) of tungsten, allowing high strength andhigh toughness to be manifested, (2) there is little chance forrecrystallization embrittlement because the material undergoes littlestructural change when heated at high temperatures due to the originalrecrystallized state, resulting in extremely little loss of strength orductility, (3) irradiation embrittlement can be greatly resolved, and(4) effects can be obtained such as a suitable decrease in yieldstrength, because the tungsten alloy crystal grain diameters grow toabout 0.05 to 10 μm when tungsten is used as the metal matrix, and atungsten alloy thus can be produced which can undergo plasticdeformation, even near room temperature.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 shows the relationship between strength and temperature for anormal metal and tungsten;

FIG. 2 shows the principle of superplastic deformation;

FIG. 3 schematically shows plastic working with the objective ofintroducing work-deformed structures with dislocations as carriers,resulting in a decrease in recrystallization temperature and anisotropy;

FIG. 4 schematically shows the GSMM step;

FIG. 5 shows the three-point bending behavior at a temperature of 400 Kin Embodiment 4 (DBTT: 310 K) and Embodiment 6 (DBTT: 420 K);

FIG. 6 shows the three-point bending behavior at 300 K in Embodiment 4;

FIG. 7 shows the X-ray diffraction pattern of Embodiment 2 (GSMMtreated) and Comparative Example 1 (not GSMM treated);

FIG. 8 is a photograph which shows the transmission electron micrographsof Comparative Example 1 and Embodiment 2;

FIG. 9 shows the X-ray diffraction patterns of Embodiment 5 (GSMMtreated) and the as-HIP prior to the GSMM treatment in Embodiment 5;

FIG. 10 is a photograph showing the transmission electron micrograph ofthe tungsten alloy of Embodiment 2.

MODE FOR CARRYING OUT THE INVENTION

The present invention is characterized in that an alloy is manufacturedby a step for degassing a raw material by heating as necessary, a stepfor subjecting the raw material that is obtained in the degassing stepto mechanical alloying (MA; also referred to as “MA step” below), a stepfor sintering the raw material powder obtained in the mechanicalalloying step using hot isostatic pressing (HIP; also referred to belowas “HIP step”), and a step for subjecting the alloy obtained in thesintering step to a recrystallization random grain boundarystrengthening treatment carried out by superplastic deformation that canmaximally utilize grain boundary sliding (also referred to below as“GSMM step,” where “GSMM” is an abbreviation for grain boundarysliding-based microstructure modification). In addition, the presentinvention is more specifically characterized in that the alloy that hasbeen manufactured by this method is a tungsten alloy. The presentinvention shall be described in greater detail below.

The raw materials that are used in the present invention will first bedescribed. The transition metal carbide that is used in the presentinvention refers to a carbide of a transition metal that is selectedfrom group IVA, VA, or VIA. Titanium carbide, zirconium carbide, niobiumcarbide, tantalum carbide, and the like are particularly preferred,because these transition metal elements rapidly diffuse and react withcarbon and tend to form carbides before the formation of brittle W₂C,and because the carbides that are formed are thermally stable. Thesetransition metal carbides of group IVA, VA, and VIA (referred to belowsimply as “transition metal carbides”) may be used individually, ormultiple carbides may be used in combinations.

The added amount of transition metal carbide with respect to the alloyis preferably 0.25 to 5 mass %. If the added amount of transition metalcarbide is less than 0.25 mass %, then the grain boundary strengtheningeffects or the migration inhibitory effects of the grain boundaries athigh temperatures will be poor, and the effect of an increase inrecrystallization temperature or the effect of inhibiting the productionof coarse crystal grains subsequent to recrystallization will be poor.There will also be insufficient remedy in low-temperature embrittlement,recrystallization embrittlement, and neutron irradiation embrittlement,as well as insufficient increase in high-temperature strength. On theother hand, the alloy will tend to have an undesirable increase inbrittleness if the added amount of transition metal carbide exceeds 5mass %.

Examples of alloy raw materials other than the transition metal carbidesinclude one or more selected from tungsten, molybdenum, vanadium,yttrium, chrome, niobium, tantalum, titanium, zirconium, hafnium, andthe like, or stainless steel, steel, and the like. However, themanufacturing method of the present invention is particularly useful forgroup VIA transition metals such as tungsten. The alloy raw materialpowder preferably has a Fischer particle diameter of 2 μm or greater.Although described in detail in the manufacturing methods presentedbelow, the reason is that high concentrations of oxygen or nitrogen inthe alloy that has been manufactured result in 1) impeding of transitionmetal carbide grain boundary precipitation/segregation which isnecessary for dramatically resolving low-temperature embrittlement,recrystallization embrittlement, and irradiation embrittlement, 2)promotion of the formation of W₂C, which is itself brittle and acts asan origin for fracture, and 3) formation of pores by oxygen and nitrogenwhich act as origins for fracture. For this reason, in order tostrengthen the weak grain boundaries in recrystallized microstructureswithin the alloy, it is essential to reduce the oxygen and nitrogencontents of the alloy raw material powder. It is preferable to carry outthe degassing step described below and to provide the raw material withthe particle diameter described above. However, if atmospheric controlis carried out strictly so as to suppress the admixture of impurities,then this requirement of 2 μm or greater may not be needed, and 1 μm orless may be sufficient.

The respective steps of the manufacturing method of the presentinvention are described below. The degassing step in which the rawmaterial powder is heated is carried out in order to decrease the finalcontent of oxygen and nitrogen impurities in the alloy. In the rawmaterial powder preparation stage, this is done in order to thoroughlyeliminate the air (in particular moisture) contained in the raw materialpowder. The degree of damage caused by the nitrogen or oxygen isdifferent depending on the metal material, and so the degassingconditions of the degassing step may be appropriately adjusted inaccordance with the metal material. With vanadium, for example, theoxygen or nitrogen will be absorbed to produce a solid solution, evenwhen heated in an ultra-high vacuum. This results in embrittlement(environmental embrittlement), and so the degassing step is carried outat a fairly low temperature or not at all. In addition, with SUS316L,there is no need to strictly carry out the step. With tungsten, on theother hand, oxygen or nitrogen that remains in the alloy precipitates orsegregates at the weak recrystallized grain boundaries as describedabove, promoting grain boundary embrittlement (recrystallizationembrittlement). Along therewith, pores are formed which act as originsof fracture. Thus, for example, when commercial tungsten powder is usedas a raw material, the raw material powder is preferably placed in acontainer at the time of preparation of the raw material powder (a boatmade of Mo or the like that is used as a powder carrier), and the rawmaterial powder is subjected to a degassing treatment at 800 to 1500° C.by evacuation to 10⁻⁴ Pa or less. However, the degassing step can beomitted, for example, by using an ultra-high purity W powdermanufactured by PLANSEE Japan Ltd. or other tungsten raw material thatalready has sufficiently low oxygen and nitrogen concentrations, the rawmaterial is sealed in an inert gas or reducing gas atmosphere (gas thathas been purified to a level at which there is negligible water contentor the like that contains the impurities) and an MA step or the like iscarried out, thereby removing the admixed oxygen and nitrogen.

Degassing is preferably carried out using a degassing time of 120 min at800° C. or greater, or a degassing time of 90 min or greater at 950° C.or greater, as a guideline. If the degassing temperature is less than800° C., then desorption of gas will be insufficient, whereas reactionswith the degassing container (boat made of Mo or the like) will tend tooccur if the temperature exceeds 1500° C., resulting in initiation ofaggregation of the raw material powder and additional undesirableeffects in subsequent steps.

With tungsten alloy, the oxygen content in the manufactured alloy is 950ppm or less, preferably 850 ppm or less, more preferably 300 ppm orless, and the nitrogen content is 60 ppm or less, preferably 50 ppm orless. If the oxygen and nitrogen contents of the alloy are below thesevalues, then the production of a dense alloy will be possible. Theoxygen or nitrogen content of the tungsten alloy at the raw materialpowder stage is about three times that of the final alloy that has beenmanufactured. Consequently, in regard to process management, it ispreferable for process management to be carried out so that oxygen iscontained at about 3000 ppm or less and nitrogen is contained at about180 ppm or less in the raw material powder upon completion of thedegassing step.

An MA step is carried out after the degassing step. With this MA step,operations extending from ball-milling of the raw material powderthrough sealing of the ball-milled powder in a capsule to produce aHIPed compact are preferably carried out in an inert gas or reducing gasatmosphere in order to prevent the admixture of oxygen or nitrogen. Ar,helium, neon, and the like are examples of inert gases, and hydrogen andthe like are examples of reducing gasses.

The MA step is a step in which high mechanical energy is imparted to theraw material alloy powder and the transition metal carbide, therebydecomposing the transition metal carbide and bringing about soliddissolution thereof in a uniform solid solution in the matrix phasealloy structure. Simultaneously, the ultrafine powder of the motherphase (alloy) is produced. This step is carried out, for example, usinga device such as a triaxial vibrating ball mill, a planetary ball mill,or an attriter. The MA treatment normally involves introducing balls andraw material powder and/or pre-alloy powder into a pot (container) androtating or vibrating this pot on a ball mill support stand, therebyimparting high mechanical energy to the raw material powder and/orpre-alloy powder. As a result, the different element species that havebeen added can be forcibly solid-dissolved, even in systems in whichsolid dissolution will not occur under the equilibrium conditions. Inaddition, extremely fine crystal grains (10 to 30 nm) can be produced ataround room temperature. In order to remove impurities from the ballsurfaces and the inner walls of the pot which are used in the MA step,the pot and balls alone may be heated under vacuum for 3 to 10 h at 150to 200° C. prior to introducing the raw material powder into the pot.

The specific treatment conditions for the MA step, such as the treatmenttime, the rotation rate, the ball material and diameter, the ratio ofthe total ball mass to total raw material powder mass, and the ratio oftotal internal container volume to total ball volume, may be determinedappropriately so that the transition metal carbide is uniformly mixed,decomposed and solid-dissolved in the alloy, so that ultrafine matrixphase metal crystals are produced, and so that effects of admixture ofcontainer and ball material into the raw material powder during the MAstep is inhibited (suppression of admixed amounts to negligible levels,or use of materials that will not affect subsequent materialcharacteristics, even if they become admixed).

The HIP step involves isostatic pressing of the MA powder produced inthe MA step using Ar gas and carrying out sintering at a comparativelylow temperature at which grain growth of the ultrafine alloy powderproduced in the MA step does not easily occur, while preventing exposureto atmosphere constituted by gas impurities that are harmful to thealloy. As a result, the transition metal carbide that is forciblysolid-dissolved during the MA step segregates and precipitates, therebypreventing grain growth of the ultrafine particles due to a pinningeffect, while also producing equiaxed ultrafine grains of alloy matrixphase in which transition metal carbide has segregated/precipitated atthe grain boundaries without strain resulting from recrystallization.Specifically, the MA powder is sealed in the inert gas or reducing gasatmosphere described above in a metal container made of soft (mild)steel, SUS, Ti, Nb, Ta, or the like. After removing the sealed gas byproducing a high vacuum (typical evacuation level of 10⁻⁴ to 10⁻⁶ Pa),the material is sintered for 1 to 5 h at 1350 to 1400° C. and 100 to1000 MPa, thereby producing an alloy having the structure describedabove. In order to remove impurities and the like on the inner walls ofthe metal container that is used in the HIP step, the metal containeralone may be heated under vacuum for 1 to 3 h at 500 to 1000° C. priorto introduction of the MA powder into the metal container.

The GSMM step is a step in which the weak grain boundaries in therecrystallized microstructure are replaced with strong transition metalcarbide heterophase interfaces or strong grain boundaries in which theconstitutive elements of the transition metal carbide have precipitatedor segregated. Consequently, when the transition metal carbideprecipitates or segregates at the grain boundaries, the grain boundarybond strength is increased, having the beneficial effect of increasingthe fracture strength and remedying embrittlement. In addition, the GSMMstep has the effect of increasing the crystal grain diameter to anappropriate size, decreasing yield strength (flow strength), andmanifesting ductility (decreasing (relaxing) the grain boundary load);removing residual gas pores that tend to act as origins for fracture (1to 3% remaining after HIP), as well as strengthening the interfaces(boundary interfaces) with different precipitates in dispersionstrengthened alloys containing precipitates (e.g., vanadium or stainlesssteel). In the present invention, the effect of grain boundary slidingat high temperatures is used in order to promote and optimize grainboundary precipitation and segregation of the transition metal carbide.FIG. 2 is a diagram showing the principle of superplastic deformation bygrain boundary sliding. Grain boundary sliding refers to crystal graindisplacement/movement in a state in which the equiaxed condition ismaintained, as indicated in FIG. 2 (2)→(3)→(4) when shear stress τ isapplied to the crystal structure in FIG. 2 (1). As a result of anextremely large number of repetitions of this type of grain boundarysliding, the transition metal carbide precipitates or segregates at thegrain boundaries, having the effect of increasing the fracture strengthat the weak recrystallized grain boundaries until it surpasses the yieldstrength (flow strength), allowing apparent alloy deformation.

However, grain boundary sliding is non-uniform deformation thatconversely promotes embrittlement along with grain boundary displacementdue to crack formation at grain boundary triple points (an example beinghigh-temperature embrittlement typically seen with copper alloys and thelike). It is thus extremely important, in the present invention, for thedeformation amount at break to be extremely large, and to employsuperplastic deformation which can maximally utilize grain boundarysliding. As stated above, grain boundary sliding is non-uniformdeformation that typically promotes embrittlement due to crack formationat grain boundary triple points in conjunction with grain boundarydisplacement. However, by carrying out superplastic deformation of thepresent invention using constant conditions described below fortemperature and strain rate (a quantity obtained by dividing the speedat which a sample piece is deformed by the size of the sample piece toconvert to strain), a relaxation (accommodation) mechanism operatingwith grain boundary sliding prevents the formation of cracks, andelongation of several hundreds of percent is produced. In order topromote and optimize transition metal carbide grain boundaryprecipitation and segregation, it is more effective to carry out GSMMfor a longer period of time than for a shorter period of time.

As stated above, superplastic deformation is a deformation mode in whichelongation of several hundreds of percent occurs due to grain boundarysliding, and equiaxed crystal grains are essentially maintained evenafter deformation. Thus, it is possible to “promote and optimizetransition metal carbide grain boundary precipitation and segregationthough relative motion or rotation of crystal grains by active grainboundary sliding over a long period of time, and to maintain anisotropic recrystallized structure with little anisotropy.” This“recrystallized random grain boundary strengthening treatment carriedout by superplastic deformation that can maximally utilize grainboundary sliding” is designated, in the present invention, by GSMM(grain boundary sliding-based microstructural modification).

FIG. 3 is a diagram that schematically shows the concept of “plasticworking performed to introduce work-deformed structures usingdislocations as a carrier, resulting in anisotropy and a decrease inrecrystallization temperature,” which has been widely used forincreasing toughness in metals including tungsten. The term“dislocation” refers to linear lattice defects, and a characteristics ofplastic working include (1) that specific crystallographic planes canperform sliding motion under small stress in specific crystallographicdirections, (2) that dislocations can multiply anew through the slipprocess, (3) that extremely strong interactions occur with areas havingan elastic strain field (e.g., the dislocations are entirely enclosed byan elastic strain field due to elastic strain that arises at around theperiphery of different atomic species with different sizes). For thisreason, when deformation progresses as a result of tensile stressapplied to a material as shown in FIG. 3(1), 3(2), slip arises in thematerial as shown in FIG. 3(3), dislocations in the material multiply(specifically, the dislocation density increases), and, as a result, thestress required for the dislocations to undergo additional slipping, inother words, the stress required for plastic deformation of the alloy,increases, and reaches the fracture strength, resulting in break asshown in FIG. 3(4). “Structures in which the dislocation density hasincreased due to plastic deformation” are referred to as work-deformedstructures or worked structures, but the increase in dislocation densityamounts to an increase in the internal strain field of the material(crystal) and produces a condition of high internal energy. Thus,materials having this high internal energy tend to release this internalenergy, and so when heat is applied (the temperature is increased), theinternal energy is released with just a slight amount of energy (slightincrease in temperature); one process for this release beingrecrystallization. In most cases, the upper limit of elongation by“plastic working performed to introduce deformation-processed structuresusing dislocations as a carrier and to produce anisotropy and a decreasein the recrystallization temperature” is approximately several tens ofpercentage points, and is considerably less than 100%, particularly withelongation.

On the other hand, with superplastic deformation occurring by grainboundary sliding, recrystallized structure having little strain ismaintained even after deformation, and the internal energy is notsubstantially increased. Thus, with the GSMM treatment of the presentinvention, although the crystal grains grow (increase by roughly afactor of ten) because the treatment is carried out at a temperaturethat is higher than the HIP temperature, grain growth leads to decreasein the total area of crystal grain boundaries and hence internal energybecause crystal grain boundaries are a high energy region.

As stated above, GSMM in the present invention is a new structurecontrol technique whereby increased durability is manifested bystrengthening the weak recrystallized grain boundaries which are a causeof grain boundary embrittlement. The plastic work referred to above issubstantially different in principle, and the post-treatment alloyfracture strength and elongation at break are also completely different.

The GSMM step, as shown in FIG. 4 involves sandwiching the alloy that isproduced in the HIP step between heat resistant, high strength ceramicsor ceramics composite plates (typically BN-SiC composite materialplates) and applying a pressure at a strain rate of 10⁻⁵ s⁻¹ to 10⁻² s⁻¹at a high temperature of 500 to 2000° C. (40 to 50% or more higher thanthe melting points of the respective alloys measured in absolutetemperature) to carry out plastic deformation due to grain boundarysliding at 60% or greater. The temperature is preferably suitablyadjusted in accordance with the melting points of the respective alloys,as stated above. For example, 1200 to 2000° C. is preferred for tungstenand molybdenum, but 1400 to 2000° C. is additionally preferred fortungsten. In addition, a temperature of 800 to 1500° C. is preferred forvanadium and SUS316. With tungsten, there are cases where the alloy willbreak during compression deformation if the temperature is less than1400° C., whereas exceeding 2000° C. is undesirable because theequipment used for industrial manufacture will increase in size. Inaddition, if the strain rate is slower than 10⁻⁵ s⁻¹, effects will beobtained, but an excessively long processing time will be required,which is industrially disadvantageous. It is undesirable for the strainrate to be greater than 10⁻² s⁻¹ due to the danger of alloy fracture.Carrying out plastic deformation of 60% or greater means that theelongation (deformation) of the test piece due to plastic deformation is60% or greater. Elongation is expressed as the elongation length of atest piece (ΔL) divided by the initial length (L), multiplied by 100 inorder to obtain a percentage. The material may instead be subjected totensile deformation, shear deformation, or the like may be used,provided that the aforementioned temperature, strain rate, and plasticdeformation can be provided.

The transition metal carbide is necessary in order for the crystalgrains of the alloy matrix phase to be maintained in fine grain sizesand in order to manifest superplastic deformation. In addition, theheterophase interface between the transition metal carbide and the alloymother phase (matrix) satisfies the Kurdjumov-Sachs azimuth(orientation) relationship, and thus high-strength hetero-phaseinterfaces are formed. When tungsten is used as the alloy raw material,90% or greater of the azimuths of the transition metal carbide presentin the tungsten alloy structure and the azimuths of the tungsten matrixare in the (Kurdjumov-Sachs) azimuth relationship: {111} W//{110}transition metal carbide <110> W//<111> transition metal carbide. If 10%or more of the transition metal carbide particles do not satisfy theKurdjumov-Sachs azimuth relationship, then it will not be possible toobtain sufficient maximum flexural strength at room temperature(approximately 1470 MPa).

In addition, with the alloy that has been manufactured by themanufacturing method of the present invention, growth occurs until thecrystal particle diameter of the tungsten alloy is about 0.05 to 10 μm.As a result, an effect is produced whereby the yield strength isdecreased to an optimal level, and a tungsten alloy can be produced thatcan undergo plastic deformation near room temperature. For tungstenalloys, when the alloy is manufactured by the manufacturing method ofthe present invention, the three-point bending ductile-brittletransition temperature (nil ductility temperature; DBTT) can bedecreased to about 500K, and thus plastic deformation is possible at orabove the ductile-brittle transition temperature.

The crystal grain diameter can be determined as the average graindiameter by using commercial image processing software (e.g., Image Pro)to carry out image processing on photographs that are typically taken bya transmission electron microscope from the center part of a samplecross section. The average grain diameter can be determined only for thetungsten matrix phase. Because averaged data could be obtained bycounting the tungsten crystal grains over a surface area ratio of 80% orgreater, statistical determinations were carried out.

For less than 20% of the surface area, counting the tungsten crystalgrains was difficult (the grains were fine and too numerous, and it wasdifficult to determine where the crystal grain boundaries were, becauseit was difficult to see the borders of the crystal grains constitutingthe crystal grain boundaries). However, if the average grain diameterfor tungsten is calculated over a region constituting 80% or more of thesurface area, then the characteristics of the various materials can beelucidated. The crystal grain diameter can be calculated as a stableaverage grain diameter by counting 300 or more tungsten crystal grainsand calculating the surface area. As necessary, the crystal graindiameter can be measured over a broad region of 80% or more of the totalfield of numerous photographs taken with a transmission electronmicroscope. As a result, 80% or more of the crystal grains that can bemeasured should be in the grain diameter range of 0.05 to 10 μm.

If the average grain diameter is less than 0.05 μm, plastic deformationwill be extremely difficult, because the yield strength will becomeextremely high, resulting in decreased work and manufacture yields,which is industrially disadvantageous. On the other hand, if the averagegrain diameter exceeds 10 μm, superplastic deformation will not readilyoccur. In order to allow plastic deformation in the vicinity of roomtemperature, it is necessary to carry out suitable adjustments so that asuitable work ratio is produced during plastic deformation (in otherwords, during GSMM treatment) that is carried out in order to increasetoughness. The temperature during the GSMM treatment may be decreased inorder to produce a smaller average grain diameter, or the temperatureduring GSMM treatment may be increased in order to produce a largeraverage grain diameter.

One point that should be noted in regard to the description of thepresent invention is that superior characteristics (e.g., fracturestrength and ductility) can be obtained in a structure in which thorough(full) recrystallization has occurred, because equiaxed crystal grainsthat are not anisotropic are produced in the metal structure. The term“equiaxed crystal grains” in the present invention means that the aspectratio (ratio of the longitudinal and transverse crystal grain lengths)is 2 or less regardless of the cross-section when the metal structure isviewed two-dimensionally.

Embodiments

With the methods for manufacturing the alloys and the manufacturedalloys of the embodiments described below, as shown in FIG. 4, simplecompressive deformation was carried out along one axis, but deformationis not restricted to simple compression, provided that superplasticdeformation allowing maximal utilization of grain boundary sliding canbe realized. Depending on the shape of the alloy product that isdesired, for example, reduction by rolling can be employed forsheet-form materials, for example.

Characterization of transition metal carbide amount required formanifesting superplasticity

Experiment 1

TiC powder with an average particle diameter of 0.7 μm (manufactured bySoekawa Chemical Co., Ltd.) was added to tungsten powder with an averageparticle diameter of 4 μm (Manufactured by A.L.M.T. Corp.) using theFischer method. The material was introduced into a molybdenum boat in ahydrogen atmosphere and was then subjected to a degassing treatment byheating for 1.5 h at 950° C. under high vacuum (<1×10⁻⁴ Pa). Next, thematerial was subjected to a mechanical alloying (MA) treatment by mixingfor 70 h at a vibration frequency of 360 cycles/min (6 Hz) in a TZM(titanium, zirconium-containing molybdenum alloy) container (pot) usinga tri-axial vibrating ball mill (TKMAC 1200, manufactured by TopologySystems). In order to characterizes the appropriate TiC powder additionrange, eight MA treatment sample runs were carried out with TiC powdercontents of 0 to 6.0 mass %.

Next, the MA-treated powder was introduced into a molybdenum boat andwas heated for 1.5 h at 950° C. under high vacuum in order to degas thehydrogen that had admixed in the TiC powder and the tungsten during theMA treatment. This degassed powder was then sealed in an HIP capsule(mild steel), and the container (capsule) was vacuum-sealed beforecarrying out a HIP treatment for 3 h at 1350° C. and 196 MPa in argongas to obtain a sintered body. The resulting sintered body is referredto as “as-HIPed compact”.

Pieces having dimensions of 0.4 mm×4 mm×16 mm were then wire-cut fromthe as-HIPed compact (parallel part length 5 mm; I-shaped flatsheet-form tensile test piece similar to the test piece shown in FIG. 1of T. Kuwabara, H. Kurishita, M. Hasegawa, Development of an Ultra-FineGrained V-1.7 mass % Y alloy Dispersed with Yttrium Compounds HavingSuperior Ductility and High Strength, Mater. Sci. Eng. A 417 (2006)16-23). The entire surface was mechanically polished with waterproofpaper (to #1500), the four edges were then chamfered, and the piece wasmounted on a tensile test fixture and subjected to high-temperaturetensile testing. The tensile fixture was a test-piece shoulder-bearing”(R part) type whereby alignment is ensured by a system in which thecompressive load on the fixture is converted to tensile load on the testpiece, allowing one-touch mounting of the test piece on the fixture.Heating of the test piece was carried out by high-frequency inductionheating using a graphite susceptor. The surface temperature of the testpiece was continually observed and recorded using a two-color radiationthermometer (Chino, model 1R-AQ). The tensile test was carried out usingan Instron model 81362 Electrically Actuated Tester at temperatures of1500° C., 1600° C., and 1700° C., an initial strain rate of 5×10⁻⁴/s(cross-head speed: 0.0025 mm/s), and an evacuation level of 5×10⁻⁴ Pa.Load and elongation (%) were measured during the tensile test. Theoxygen and nitrogen concentrations in the sample were measured byinfrared absorption on a LECO-TC600 device using a thermal conductivitymethod. With all of the samples, the oxygen concentration was 850 ppm orless, and the nitrogen concentration was 50 ppm or less. The results areshown in Table 1. In the table, >160 denotes that no break occurred,even at a deformation of 160%.

TABLE 1 1500° C. 1600° C. 1700° C. TiC content Elongation ElongationElongation Sample No. Mass % (%) (%) (%) 1 0 2 5 5 2 0.15 3 10 30 3 0.2570 105 >160 4 0.5 >160 >160 >160 5 1.1 >160 >160 >160 61.5 >160 >160 >160 7 5 >70 >100 >160 8 6 10 30 60

Experiment 2

The same test as in Experiment 1 was carried out, with the exceptionthat the hydrogen in Experiment 1 was changed to argon, and nine sampleswere used in which the TiC content was varied. The results are shown inTable 2

TABLE 2 1500° C. 1600° C. 1700° C. TiC content Elongation ElongationElongation Sample No. Mass % (%) (%) (%) 9 0 2 5 5 10 0.25 3 7 7 11 0.530 40 60 12 0.7 50 70 >160 13 0.8 70 110 >160 14 1.1 120 >160 >160 151.5 70 >160 >160 16 5 50 70 >160 17 6 10 30 60

Experiment 1 and Experiment 2 above show that the TiC amount requiredfor manifestation of superplasticity at 1600 to 1700° C. (elongation atbreak:100% or greater) is 0.25 to 5 mass % with the as-HIPed compactsproduced from powder that was MA-treated in hydrogen atmosphere, and 0.7to 5 mass % with those produced using an argon atmosphere. If the TiCamount is below these ranges, then weak grain boundaries occur in greatnumbers among the grain boundaries of the tungsten phase, and there arefew grains of a second phase that inhibit grain boundary movement. Forthis reason, grain growth is rapid in the tungsten phase, resulting inthe production of large-size crystal grains. The TiC phase is essentialfor rotation and movement of crystal grains during grain boundarysliding and for maintaining fine equiaxed crystal grains that arerequired for superplastic deformation. For this reason, if the amount ofTiC phase is small, then when grain boundary sliding non-uniformdeformation arises in high-temperature tensile testing, grain boundarycracks will form and grow, and the elongation at break will be low.

Conversely, if the amount of TiC exceeds these ranges, then the contactfrequency between TiC phases will increase, and the proportion ofTiC/TiC interfaces will increase. In comparison to the tungsten matrixphase, the TiC phase has low plastic deformation capability, and it isthought that TiC/TiC interfaces also do not readily slide. Consequently,overloading of the harmony of the tungsten phase occurs in relation tocontinuous grain boundary sliding of the tungsten grains, and grainboundary (interfacial) cracking arises, thereby also decreasing theelongation at break.

Experiment 3

An experiment was carried out in the same manner as in Experiment 1,with the exception that the material was changed to titanium carbide ofExperiment 1, and zirconium carbide, niobium carbide, tungsten carbide,or mixtures thereof were added while varying the content thereof. Inaddition, high-temperature tensile testing was carried out only at 1600°C. The results are shown in Table 3.

TABLE 3 1600° C. Carbide content in tungsten Elongation (%) ZrC 0.3 mass% 120 ZrC 4.7 mass % >160 NbC 0.32 mass % 130 NbC 4.5 mass % >160 TaC0.28 mass % 130 TaC 3.3 mass % >160 TaC 5.0 mass % >160 ZrC 0.3 mass % +TaC 2 mass % >160 NbC 0.3 mass % + TiC 2 mass % >160 TiC 1 mass % + TaC2 mass % >160 TaC 0.1 mass % 2 TiC 0.08 mass % + TaC 0.03 5 mass % ZrC 6mass % Could not be performed

As is clear from Table 3, even when transition metal carbides other thanTiC were added to the alloy at about 0.25 to 5 mass %, it was clear thatthere was an improvement in ductile characteristics.

Embodiment 1

An as-HIPed compact material was produced in the same manner as withsample no. 5 of Experiment 1, with the exception that the degassingconditions involved heating for 1.5 h at 1050° C. under high vacuum(1×10⁻⁴ Pa). Next, the as-HIPed material that had been produced waswire-cut to produce a sintered body with a diameter of about 9 to 10 mmand a height of about 20 mm. In order to strengthen the weak randomgrain boundaries by superplastic deformation that can maximally utilizegrain boundary sliding, a disk shape material was produced bycompression deformation to a thickness of about 3.5 mm (diameter ofabout 21 to 23 mm) at a temperature of 1650° C. and a strain rate of 0.5to 2×10⁻⁴ s⁻¹ (superplastic behavior tends to more readily occur withslower strain rates, and so a rate was selected at which experimentswere most easily carried out, while considering response (increase inflow stress) exhibited by the material as a result of incrementalincreases in strain rates). Heating of the sintered body was carried outby high-frequency induction heating under evacuation using a graphitesusceptor. An Instron model R1362 Electrically Actuated Tester was usedfor high-temperature compressive deformation. A piece with dimensions of1 mm×1 mm×20 mm was cut perpendicularly in the compressive directionfrom this disk shape material, and the surfaces and edges were polishedwith water-proof emery paper out to #1500, thereby producing a bendingtest piece. The oxygen concentration was 40 ppm and the nitrogenconcentration was 30 ppm in the test piece, as measured by infraredabsorption on a LECO-TC600 device using a thermal conductivity method.Next, the test piece was subjected to three-point bend testing at atemperature range of room temperature to 600° C. with a cross-head speedof 0.001 mm/s in an atmosphere produced by a flow of high-purity Arcontaining 4% H₂. The three-point bend test was carried out using aServopulser EHF-2 model fatigue testing machine, manufactured byShimadzu Corporation (load capacity of 5 ton), connecting a span ±2.5 mmLVDT (Linear Variable Differential Transformer) to the actuator head andattaching a shear-type load sensor with a load capacity of 5 kN directlybelow a load cell with a capacity of 5 ton. Control of testing wascarried out using a static test application program. An infrared heatingfurnace (ULVAC-RIKO, Inc.) was used for heating the test pieces, andmeasurement of the test piece temperature and atmosphere (at a locationseparated by a few millimeters from the test piece) was carried out inadvance on a dummy test piece having a spot-welded thermocouple. Inactual testing, the temperature of the atmosphere was controlled andmeasured. Flexural strength was measured at room temperature, and theminimum value of the measured averages of five bending test pieces wastaken as the minimum flexural strength, whereas the maximum value wastaken as the maximum flexural strength. The DBTT was determined byrecording the variation of measurements on the plastic strain atrespective temperatures while increasing the testing temperature byroughly 50 increments starting from room temperature. The temperaturefound by extrapolating to a plastic strain of zero using a linearapproximation was taken as the DBTT. In determining a single DBTT, it isnecessary to measure the plastic strain while varying the testingtemperature, and measurements were carried out by preparing three tofive bar shape test pieces having the same impurity concentrations andstructures.

Embodiments 2 to 7

The degassing conditions involved heating for 1.5 h at 950° C. inEmbodiment 2, heating for 1 h at 950° C. in Embodiment 3, heating for 1h at 900° C. in Embodiment 4, heating for 1.5 h at 850° C. in Embodiment5, heating for 1 h at 850° C. in Embodiment 6, and heating for 1 h at800° C. in Embodiment 7. With the exception that the oxygen amount andnitrogen amount in the tungsten alloy were changed, test pieces wereproduced out using the same procedure as in Embodiment 1. The oxygenamount, nitrogen amount, minimum flexural strength and maximum flexuralstrength at room temperature, and DBTT were measured.

Comparative Example 1

Test pieces were prepared from as-HIPed compacts without carrying out acompressive deformation treatment for GSMM, and measurements wereperformed using the same procedure as in Embodiment 2.

Comparative Example 2

With the exception that the TiC content was changed to 1.1 mass % andthe degassing treatment was not carried out, test pieces were preparedby the same procedure as in Embodiment 1, and measurements wereperformed.

The results of measurements in Embodiments 1 to 7 and ComparativeExamples 1 and 2 are shown in Table 4.

TABLE 4 Minimum Maximum Oxygen Nitrogen flexure flexure amount amountTiC strength strength DBTT (ppm) (ppm) (%) (MPa) (MPa) (K) Embodiments 140 30 1.1 2800 3200 210 2 160 30 1.1 2700 2800 230 3 230 40 1.1 26902940 240 4 610 40 1.1 1840 2380 310 5 850 50 1.1 1450 1620 330 6 870 1401.1 1240 1500 420 7 950 60 1.1 1340 1470 500 Comparative 1 160 30 1.11610 2160  850* examples 2 2120 180 1.1 1000 1260 ≧630   

As is clear from Table 4, the tungsten alloy flexural strength increasedas the concentrations of oxygen and nitrogen decreased when the materialwas subjected to a GSMM treatment. In addition, it was clear that theDBTT decreased dramatically when the as-HIPed material was subjected tothe GSMM treatment, and ductility was obtained even at low temperatures.

Three-Point Bending Testing

FIG. 5 shows the three-point bending behavior at a temperature of 400Kfor Experiment 4 (DBTT: 310 K) and Experiment 6 (DBTT: 420K). FIG. 6shows the three-point bending behavior for Experiment 4 at 300K. As isclear from FIGS. 5 and 6, break (fracture) occurred withoutmanifestation of ductility at lower temperatures than the DBTTtemperatures of the resulting alloys, and thus it was clear that theamounts of oxygen and nitrogen must be decreased in addition to carryingout the GSMM treatment (by compression).

X-Ray Diffraction Pattern Measurement

FIG. 7 compares the X-ray diffraction patterns for Experiment 2 (GSMMtreatment performed) and Comparative Example 1 (GSMM treatment notperformed). From a comparison of the two, a large difference inintensity was seen with the TiC peak, indicating that TiC precipitationprogressed during the GSMM treatment. This was confirmed by transmissionelectron microscopy.

Transmission Electron Micrographs

FIG. 8(1) is a transmission electron micrograph of Comparative Example1, and FIG. 8(2) is a transmission electron micrograph of Embodiment 2.FIG. 8(3) is an enlarged view of the portion indicated by “←” in FIG.8(2). As is clear from the micrographs, the tungsten alloy that had beensubjected to the GSMM treatment was confirmed to have experienced TiCgrain boundary precipitation in the alloy. It was simultaneouslyconfirmed that the TiC constituent elements had undergone soliddissolution and segregated at the grain boundaries.

X-Ray Diffraction Analyses

FIG. 9 compares the X-ray diffraction patterns of Embodiment 5 (GSMMtreated) and the as-HIPed material prior to GSMM treatment in Embodiment5. It is clear that TiC precipitation similarly progressed as a resultof GSMM treatment, but that the ductility-impeding (i.e., easilyfractured) carbide W₂C, had formed, in contrast to FIG. 7. It is thoughtthat this material is produced as a result of an increase in oxygenlevel, which provides oxygen distribution between the TiC and tungstenand results in a reaction between some of the carbon that hasdissociated from the TiC with the surrounding tungsten.

Confirmation of Equiaxed Recrystallized Grains

With the tungsten alloy of Embodiment 2, a thin film having a smallperforation was formed at the center by electrolytic polishing (TenuPol)at a thickness of about 50 μm and a diameter of 3 mm, whereupon thematerial was observed with a transmission electron microscope (JEOL2000) at an acceleration voltage of 200 kV. FIG. 10(1) shows theobservation direction of the transmission electron microscope, and FIG.10(2) shows a micrograph of the sample as observed from above(specifically, from a direction parallel to the compression direction).FIG. 10(3) is a micrograph of the sample as observed from the side (froma direction perpendicular to the compression direction). In all cases,clear images were obtained with an observational magnification ofroughly 10,000×. As is clear from FIG. 10, the crystal grains wereequiaxed grains, with crystal grain aspect ratios in the range of 1 to2.

In addition, upon observing the thin film under diffraction conditionsand at additionally high resolution, almost no dislocations wereobserved in the crystal grains, and the number of dislocations with theobserved crystal grains was extremely low, at about 1 to 3, in manycases. From the results of observations described above, it is clearthat the structure of the articles of the present invention is arecrystallized structure. Dislocations are present at 1000 or greater inthe crystal grains of tungsten that has not been recrystallized,including work-deformed structure. In contrast, it was found that thethere are 50 or less dislocations in the crystal grains ofrecrystallized tungsten. It became clear that, if this condition issatisfied, then the characteristics of unstrained tungsten crystalgrains are exhibited.

In addition, an unstrained condition was confirmed based on XRDmeasurements using a Rigaku RAD II-B device. From the results of XRDmeasurements, although the effect of fine crystal grains was obtained,the diffraction width of the diffraction peaks increased with increasingstrain in the non-recrystallized state. For example, by investigatingXRD results obtained by measurements on stress-relieved commercialtungsten material and tungsten alloy of Embodiment 2 under conditions of40 kV and 30 mA using a Cu target and a 1° slit, it was clear that ifthe full width at half maximum exceeds 3° in (220) diffraction oftungsten with a lattice constant of 0.11188 nm, strain remains and thematerial does not have a recrystallized structure (commerciallyavailable pure tungsten subjected to stress relief treatment), whereasthe material has a recrystallized structure without strain if the fullwidth at half maximum is 3° or less.

Confirmation of Grain Diameter

The tungsten alloys produced in Embodiments 1 to 7 were used, thin filmshaving a small perforation was formed at the center by electrolyticpolishing (TenuPol) at a thickness of about 50 μm and a diameter of 3mm, whereupon the materials were observed with a transmission electronmicroscope (JEOL 2000) at an acceleration voltage of 200 kV. In thetransmission electron micrographs of all of the embodiments, the crystalgrain diameters could be measured for 80% or more of the entire field ofthe micrographs, and it was confirmed that 80% or more of the crystalgrains that were measured were in the grain diameter range of 0.05 to 10μm.

Confirmation of Carbide Orientation and Tungsten Matrix Orientation inTungsten Alloy Structure

In contrast that the case described in “Confirmation of grain diameter”above, bright-field images, dark-field images, and selected areadiffraction patterns at magnifications of several hundred thousand weretaken in numerous fields including a single carbide grain within atungsten matrix phase. The orientation relationship between the carbideand the tungsten matrix phase were thus analyzed. As a result, in thetransmission electron micrographs from all of the embodiments, it wasconfirmed that the carbide present in the tungsten alloy structure andthe tungsten matrix satisfied the Kurdjumov-Sachs orientationrelationship: {111}W//{110} transition metal carbide <110> W//<111>transition metal carbide.

Embodiment 8

After weighing and blending raw material powders at avanadium:yttrium:tungsten:TiC mass ratio of 89.8:1.4:8.0:0.8, thematerial was placed in a Mo boat, and a degassing treatment was carriedout for 1 h at 200° C. Next, a container and balls used for MA treatment(material: TZM (Mo-0.5 Ti-0.1 Zr)) was baked for 10 h at 150 to 200° C.under high vacuum, whereupon the blended raw material powder wasintroduced into the container along with the balls, and an MA treatmentwas carried out using a tri-axial vibrating ball mill for 70 h in apurified hydrogen atmosphere. In order to eliminate the hydrogen thathad been admixed from the atmosphere during MA, a dehydrogenationtreatment was carried out for 1 h at 600° C. under a vacuum of 1×10⁻⁴ Paor less. Subsequently, the MA-treated vanadium alloy powder was sealedin a hydrogen atmosphere in an HIP capsule (soft (mild) steel) that hadbeen degassed by heating under vacuum at 900° C., and, while degassingunder vacuum at room temperature, the HIP capsule was vacuum-sealedunder a high vacuum (2×10⁻⁵ Pa). Consequently, the HIP capsule interiorwas in a highly evacuated tightly-sealed condition. This material wasthen subjected to an HIP treatment for 3 h at 196 MPa and 1000° C. inargon gas to produce a sintered body with a relative density of 99.5% ormore. Tensile test pieces were then cut out from the sintered body inthe same manner as in Embodiment 1, and superplastic deformationallowing maximal utilization of grain boundary sliding was employed. Inorder to strengthen weak portions, such as interface boundaries betweenprecipitates and vanadium matrix phase that generally work as crackinitiation sites, a GSMM treatment was carried out at temperature of1300° C. and a strain rate of 0.5 to 2×10⁻⁴ s⁻¹. The resulting testpiece was subjected to tensile testing under conditions of roomtemperature and an initial strain rate of 1×10⁻³/s using a ServopulserEHF-2 model device manufactured by Shimadzu Corporation. The yieldstrength (=0.2% proof strength), tensile strength, uniform elongationand elongation at break (total elongation) were measured.

Comparative Example 3

Production of test pieces and measurements were carried out using thesame procedure as in Embodiment 8, with the exception at a GSMMtreatment was not carried out.

Embodiment 9

SUS316L (316L stainless steel alloy powder supplied by Höganäs AB) andTiC were used at a SUS316L:TiC mass ratio of 98:2 as alloy raw materialpowder. A degassing treatment was carried out for 1.5 h at 450° C. Afteran MA treatment, a dehydrogenation treatment was carried out for 1.5 hat 450° C. The heating temperature while sealed under vacuum in the HIPcapsule was 750° C., and a HIP treatment was carried out for 3 h at 850to 900° C. In addition, a GSM treatment was carried out at 950° C. Withthese exceptions test pieces were produced and measurements were carriedout in the same manner as in Embodiment 8.

Comparative Example 4

Production of test pieces and measurements were carried out using thesame procedure as in Embodiment 9, with the exception that a GSMMtreatment was not used.

The results of measurement in Embodiments 8 to 9 and in ComparativeExamples 3 and 4 are shown in Table 5.

TABLE 5 0.2% Proof Tensile Uniform Elongation strength strengthelongation at break (GPa) (GPa) (%) (%) Embodiment 8 0.66 0.7 14 20 90.84 1.13 20 23 Comparative 3 0.71 0.72 5 9 Example 4 0.68 0.95 10 10

As is clear from Table 5, with vanadium alloys and stainless steelalloys that were subjected to the GSMM treatment, uniform elongation andelongation at break were at least doubled, and it was clear that theGSMM treatment can improve the ductility characteristics of varioustypes of metals or alloys, including tungsten.

INDUSTRIAL APPLICABILITY

Because low-temperature embrittlement, recrystallization embrittlement,and irradiation embrittlement of alloys, particularly tungsten, can bedramatically remedied by subjecting the alloy to a GSMM treatment, newavenues for the utilization of alloys, particularly tungsten, areexpected in regard to use in extreme environments involving exposure tosevere thermal loads, such as with high-temperature structuralmaterials, molybdenum substitute materials, plasma-facing materials forinternational thermonuclear experimental reactor (ITER),high-temperature test fixtures, and solid rotating targets forspallation neutron sources.

1. A method for manufacturing an alloy, characterized by having a stepfor mechanical alloying of a metal raw material and at least oneselected from carbides of group IVA, VA, or VIA transition metals, astep for sintering the raw material powder obtained in said mechanicalalloying step using hot isostatic pressing, and a step for subjectingthe alloy obtained in said sintering step to grain boundary slidingbased plastic deformation of 60% or greater at 500 to 2000° C. and astrain rate of 10⁻⁵ s⁻¹ to 10⁻² s⁻¹.
 2. The method for manufacturing analloy according to claim 1, characterized by having a step in which saidtransition metal carbide and the metal raw material are degassed byheating prior to said mechanical alloying step.
 3. A tungsten alloycomprising 0.25 to 5 mass % of at least one type selected from carbidesof a group IVA, VA, or VIA transition metals, the tungsten alloycharacterized in that the oxygen content is 950 ppm by mass or less, thenitrogen content is 60 ppm by mass or less, 80% or more of an observedcross sectional area in the tungsten phase is recrystallized equiaxedgrains with grain diameters of 0.05 to 10 μm, the ductile-brittletransition temperature determined by three-point flexure is 500K orless, and plastic deformation due to grain boundary sliding is possibleat or above this temperature.
 4. The tungsten alloy according to claim3, characterized in that 90% or greater of the azimuths of the carbidepresent in the tungsten alloy structure and the azimuths of the tungstenmatrix are in the (Kurdjumov-Sachs) azimuth relationship: {111} W//{110}transition metal carbide <110> W//<111> transition metal carbide.
 5. Thetungsten alloy according to claim 3, characterized in that the fullwidth at half maximum for reflection of the (220) diffraction planes is3° or less as determined by X-ray diffraction, or that there are 50 orfewer dislocations within crystal grains as determined by transmissionelectron microscopy.
 6. The tungsten alloy according to claim 3,characterized in that the maximum bend strength determined bythree-point flexure is 1470 MPa or greater.
 7. The alloy that ismanufactured by the manufacturing method according to claim
 1. 8. Thetungsten alloy according to claim 4, characterized in that the fullwidth at half maximum for reflection of the (220) diffraction planes is3° or less as determined by X-ray diffraction, or that there are 50 orfewer dislocations within crystal grains as determined by transmissionelectron microscopy.
 9. The tungsten alloy according to claim 4,characterized in that the maximum bend strength determined bythree-point flexure is 1470 MPa or greater.
 10. The tungsten alloyaccording to claim 5, characterized in that the maximum bend strengthdetermined by three-point flexure is 1470 MPa or greater.
 11. The alloythat is manufactured by the manufacturing method according to claim 2.